Probing the effect of ethylene carbonate on optimizing the halogen-free electrolyte performance for Mg sulfur batteries

Magnesium metal batteries attract great attention for their high volumetric capacity and safety as a post-lithium choice. The strategy of adding organic plasticizer may bring new insights into designing halogen-free electrolytes for the further development of magnesium–sulfur batteries. The high charge density of Mg2+ results in a high desolvation barrier and low interfacial Mg2+ transfer kinetics due to the strong coulombic interactions of Mg2+ ions with anions and solvent molecules. In this study, we test the effect of the stoichiometric ratio of ethylene carbonate (EC) as an organic additive on the electrochemical performance of halogen-free electrolyte (HFE) based on Mg(NO3)2 in acetonitrile (ACN) and tetraethylene glycol dimethyl ether (G4). Through various characterization methods, the introduction of EC perturbs the bonding scheme of the HFE electrolyte, enhances the ionic conductivity, reduces the relaxation time, and forms a resistive solid electrolyte interphase (SEI). The assembled Mg–S full cell using modified HFE (HFE_EC) delivers initial specific capacities of 900 m Ag−1 with a cycle life of up to 10 cycles in the case of activating the cell with electrochemical conditioning. This study sheds light on the interplay of EC and the interfacial kinetics in Mg batteries and opens a door for designing novel magnesium electrolytes.


Introduction
The transition from fossil fuel to green fuel depends on the ability to bridge the gap of renewable energy intermittency through sustainable energy storage systems. Mg-S batteries are an attractive option due to their low cost, environmental sustainability, and abundant magnesium and sulfur in the Earth's crust. In addition, Mg-S batteries offer competitive theoretical energy density (1684 W h kg −1 and 3286 W h L −1 ) and safety (dendrite less) over their Li counterparts. However, there are technical challenges like the poor cyclability, the shuttle effect of polysuldes, and the highly resistive passive anode/electrolyte interface that must be overcome before the practical transition from the lab to the market. 1 BaTiO 3 possesses a piezoelectric property, generating electrostatic charges on the surface under electrochemical stress. This property gives BaTiO 3 the ability to immobilize the polar poly-sulde molecules via chemical adsorption, reducing polysulde shuttling and improving the cycle life of Mg-S cells. 2,3,4 The electrolyte controls the electrochemical activity, safety stability, and thermodynamics of the batteries. Therefore, designing an efficient magnesium electrolyte compatible with sulfur's electrophilic nature will be essential to achieving a practical Mg-S battery. In this regard, non-nucleophilic electrolytes were designed, including hexamethyldisilazide-Br, 5 non-nucleophilic Mg(TFSI) 2 -MgCl 2 -1,2-dimethoxyethane/tetrahydrofuran, 6 magnesium tetrakis(hexauoroisopropyloxy)borate, 7 magnesium peruorinated tert-butoxide (Mg(pb) 2 ), 8 Mg[B(Ohp) 4 ] 2 . 9 Nevertheless, the common factors of the most weakly coordinating Mg salts are the low stability, high cost, and complex synthesis process, that impede their practical applications. In our previous research, 10 we reported a simple halogen-free electrolyte (HFE) 0.69 M Mg(NO 3 ) 2 $6H 2 O in ACN : G4 (∼2 : 1) with ionic conductivity of ∼10 −4 S cm −1 at 313 K, dendrite-free, and translucent Mg deposits. However, much work need to be dedicated to lowering the Mg 2+ desolvation energy and reducing the Mg stripping/plating overpotential. Ethylene carbonate (EC) organic cosolvent is commonly used to tailor the interfacial kinetics and the transport properties of the electrolyte and played a vital role in the commercialization of lithium-ion batteries. The high dielectric constant of EC grants it the ability to dissociate metal salt and engineer the solid anode/ electrolyte interphase. 11,12 Herein, we study for the rst time the benecial effects of the interplay between HFE and EC additive and how it affects the Mg 2+ desolvation energy. The study was revealed by UV-vis spectroscopy, Fourier transforms infrared (ATR-FTIR) spectroscopy, and electrochemical impedance spectroscopy (EIS). Galvanostatic cycling was conducted on the Mg/Mg symmetric battery with the HFE_EC electrolyte to test the change of the overpotential with introduction of EC.
The nature of Mg/HFE_EC interface was probed using FTIR, Xray diffraction (XRD), and SEM_EDS techniques. Mg-S cells with HFE and HFE_EC electrolytes were assembled and tested electrochemically, and the sulfur cathode was retrieved from the disassembled cells to examine the morphology and structure of the S cathodes at different electrochemical states using SEM, EDX, XRD and EIS. The study conrms the impact of EC additive to regulate the electrochemical performance of the HFE via modulating Mg 2+ desolvation barrier but it results in undesirable SEI which causes high Mg stripping/plating overpotential.

Experimental work
The halogen-free electrolyte (HFE) based on 0.69 M Mg(NO 3 ) 2 -$ 6 H 2 O dissolved in ACN : G4 (∼2 : 1) was synthesized according to the published work. 10 Five bottles of HFE were prepared and x M of EC, (x = 0, 0.05, 0.1, 0.15, 0.2) was added to each one, which were denoted as HFE, HFE_X 1 , HFE_X 2 , HFE_X 3 , and HFE_X 4 respectively. ATR-FTIR measurements were recorded in the wavenumber ranging from 400 to 4000 cm −1 using ALPHA II Bruker spectrometer. UV-vis spectral data were collected using Edinburgh DS5 Dual Beam UV-vis spectrophotometer. The S cathode was prepared by grinding 75 wt% of sulfur (S, Alfa Aesar 99%) with 20 wt% graphene nanoplatelets (GNPs, Grade M, XG Science), and 5 wt% BTO (Alfa Aesar 99%). The resultant mixture was placed in the microwave for 10 seconds and then subjected to a ball mill for 30 hours to obtain the S_GNP_BTO active cathodic material. 0.75 g S_GNP_BTO : 0.1 g Super C P : 0.15 g polyvinylidene uoride was dissolved in N-methyl-2-pyrrolidinone using a magnetic stirrer. The resultant viscous slurry was spread on Al foil with a thickness of 100 mm using HOHSEN MC-20 Mini-Coater and le in the oven for two hours at 100°C. Aer that, the cathode was cut with a diameter of 14 mm and then permanently kept at a temperature of 65°C. Electrochemical impedance spectroscopy (EIS) and electrochemical measurements of the electrolyte were tested using CHI604E electrochemical workstation. The collected impedance data were analyzed by ZView soware. The galvanostatic test was conducted between 0.25 and 2.5 V using a NEWARE BTS4000. XRD pattern was recorded on Rigaku MiniFlex 600 diffractometer, while SEM images were captured using Jeol JMS-700 EDS electron microscope.

Results and discussions
The functional group's evolution and possible molecular interaction of HFE_ x EC electrolytes can be identied using the FTIR spectroscopic analysis. Fig. 1a-c shows the ATR-FTIR spectra of HFE_ x EC electrolytes in the 400-4000 cm −1 range. The spectra of the pristine sample exhibit spectral lines that match well with the bands of HFE that have been reported in our previous study. 13 Aer EC addition to HFE, there are two bands at 1770 and 1796 cm −1 grew up with increasing the content of EC that match well with the C]O stretching mode of pure EC. 14 Furthermore, the introduction of the EC in the structural skeleton of the HFE results in additional small new peak characteristic of -C^N/Mg 2+ at wavenumber 2362 cm −1 , which indicates the role of EC in engineering electrolyte solvation structure via supporting weak solvent-cation interaction, Fig. 1 (a, b and c) FTIR spectra, (d) (ahn) 0.5 vs. hn, (e) n vs. l, (f) k vs. l, HFE, 5%, 10%, 15% and 20%; of HFE_ x EC electrolytes.
promoting the rapid desolvation of Mg 2+ ions. 15 The UV-visible absorption spectra are a powerful tool to analyze the band structure of the energy materials. Fig. 1d-f shows the UV-visible absorption spectrum of HFE_ x EC electrolytes in the 200-1100 nm range. The band gap energy E g was estimated using Tauc's relation: 16 Fig. 1d displays the direct electronic transition prole of hn versus (ahn) 0.5 for HFE_ x EC electrolytes. The E g estimated from the extrapolation of the linear region of the plot on the x-axis (ahn = 0) gives the value of the optical bandgap, E g . The bandgap was found to be around 3.43 eV of the HFE_ x EC electrolytes; hence HFE_ x EC shows the difficulty in losing electrons (non-nucleophilic), which makes it a compatible electrolyte with the electrophilic nature of the S 8 .
versus the wavelength (l) of HFE_ x EC electrolytes. The value of k shows a high dispersion with increasing wavelength in the UV period (200-340 nm) and is followed by a little increase in the visible region. Also, it can be noticed that the value of k relatively increases with increasing EC content. The refractive index (n) is an optical property that measures the change in the velocity of light inside the electrolyte, which can be calculated using the T)], A and T are the absorption and transmission coefficient, respectively. Fig. 1f shows the spectra of refractive index versus the wavelength (l) of the relation between the refractive index (n) of HFE_ x EC electrolytes. The prole shows increasing the ability of the HFE_ x EC mediums to bend the path of the electromagnetic waves with the transition from UV to the visible regions and increasing EC content. This indicates EC's role in changing the optical density and electronic structures of HFE. Fig. 2a shows the Nyquist plots of symmetric stainless steal (SS)//electrolyte//SS coin cell at 303 K. The spectra show an incomplete semicircle at the high-frequency region representing the bulk resistance R b (due to migration of ions) and bulk capacitance (due to immobile species), whereas the lowfrequency spike is due to the interfacial polarization effect. 19 The bulk conductivity was calculated by tting the semicircle curves using ZView soware and substituting in the following  understood in the light of the low concentration of EC was not enough to dissociate the ions. Fig. 2c shows a Cole-Cole plot of the real dielectric constant 3 ′ versus imaginary dielectric loss 3 ′′ of HFE_ x EC electrolytes at 303 K, and each plot appears as a semicircle and intersects the 3 ′ -axis. Relaxation time s was calculated using the relation su max = 1, where u max is the angular frequency of the maximum 3 ′′ . Fig. 2d shows that the relaxation time decreases exponentially with the temperature increase for all concentrations of EC due to the drop in the viscosity. Fig. 2e and f Fig. 2g. A small oxidation peak at 3 V emerged in the plasticized electrolyte while the main oxidation potential onset of the plasticized electrolyte shis to higher values >4 V compared with the pristine HFE, which suggests the role of EC in perturbing the bonding scheme of the pristine matrix. The evolution of the Mg electrolyte/electrode interface is systematically examined at rest using EIS. Fig. 3a and b show the Nyquist plots of Mg/HFE/Mg and Mg/HFE_X 4 /Mg cells with different rest time at open circuit voltage. The diameter of the semicircle decreases with an increase in the rest time, then decreases with a further increasing the rest time, and eventually stays at a stable value implying a stable interface between the electrolyte and the Mg-metal electrode. Fig. 3c shows Nyquist plots of symmetric Mg/HFE/Mg and Mg/HFE_X 4 /Mg cells before and aer cycling. The plots were tted using ZView soware with the equivalent circuit shown as inset Fig. 3 Fig. 5a. The CV response shows that the cell with HFE_X 4 has an observable anodic peak at ∼2 V and a higher cathodic/anodic current compared with the cell with HFE. However, this behavior conrms that electrolytes decomposed to an unstable SEI. EIS is regarded as one of the most powerful tools to analyze the cathode kinetic performance. The Nyquist plots of Mg/HFE/S and Mg/HFE_X 4 /S cells before and aer short cycling in the frequency range between 1 Hz and 1 MHz are shown in Fig. 5b. The impedance spectra display intercept of Z ′ at high frequency, semicircle at high-range frequency, semicircle at middle-range frequency, and a sloping line at low frequencies assigned to the bulk resistance of the electrolyte, diffusion, and migration of Mg 2+ through the solid electrolyte interphase on the surface of the electrode (R SEI ), charge transfer resistance (R ct ) and to the Warburg impedance (represents magnesium-ion diffusion through the sulfur electrode), respectively. EIS parameters were calculated via tting the equivalent circuit (Randell circuit model) using ZView soware and the obtained parameters are displayed in Table 2.
The diffusion coefficient of Mg 2+ (DMg 2+ ) was calculated by , from linear tting the of the relation Z ′ = R s + R ct + su −0.5 (u −0.5 vs. Z ′ ), the Warburg factor s was calculated from the slope at low frequencies, Fig. 5c. The estimated values of D Mg 2+ for all the cells are around  6 × 10 −15 cm 2 s −1 . Fig. 5d shows the galvanostatic curves with a current density of 0.02 mA cm −2 for Mg/HFE/S and Mg/ HFE_X 4 /S cells. The EC-containing electrolyte delivered a high discharge/charge capacity of ∼3000/1450 mA h g −1 over the theoretical capacity of the S electrode (∼1673 mA h g −1 ), while the pristine electrolyte (HFE) delivered reasonable discharge/ charge below the theoretical capacity of sulfur with short cycle life. Thus, we can conclude that the two cells' major discharge/ charge capacity value is due to the electrolyte decomposition. Aiming to improve the cycle life of the developed electrolyte in the Mg-S cells, galvanostatic cycling was considered as a conditioning process to scavenge the active contaminants and partially remove the passivation from the surface of the anode. Fig. 5e shows the voltage vs. specic capacity with a current density of 0.02 mA cm −2 of Mg/HFE_X 4 /S cell aer preactivating using the electrochemical conditioning processes. The activated cell delivered an initial discharge/charge capacity of ∼875/355 mA h g −1 , and the cycle life extended to 10 cycles with coulombic efficiency >100%, suggesting that the electrolyte decomposition is still the major factor that controls the cycle life even with the electrochemical conditioning processes. The XRD was conducted to probe the crystal structure changes of the sulfur cathode from pristine / discharge / recharge states, Fig. 6a shows the XRD pattern of the sulfur, BTO powder, and sulfur cathode at different electrochemical states. XRD pattern of S 0 (pristine cathode) shows the main peak locations at 23°, 25.74°, 28.58°, which matches the 222, 026, and 313 reection signals of S 8 with orthorhombic S lattice (DB Card No.: 9011362, Fddd: 2), conrming that sulfur is the main constituent of the cathode. 24 Aer discharge, all peaks showed intense broadening indicating the formation of an amorphous electrolyte/cathode interface and the reaction of sulfur atoms with Mg 2+ . Fig. 6b shows the zoomed-in 2q = 23°in the pristine state. The peak shied to a low angle aer discharge while it showed further where b is the full width at half maximum (FWHM), q is the Bragg's angle, k is the shape factor (0.94), and l is the X-ray wavelength. Fig. 6c reveals Williamson-Hall diagram: b cos(q) versus 4 sin(q) of S 0 , S 1, and S 2 , the slope of the line and the y-intercept represent the strain and crystallite size, respectively. 25 The lattice strain values and crystallite size are listed in Table 3. The microstrain/crystallite size increases/ decreases in the discharge state while returns to the original values upon aer recharge state, which conrms that S cathode is under high volumetric expansion/contraction during the insertion/extraction of Mg 2+ within the sulfur skeleton. SEM images and EDS results of the S 0 , S 1 , and S 2 composite are shown in Fig. 6d-g. The pristine S 0 electrode shows an inhomogeneous distribution of sulfur deposits on the GNPs network surface. Aer magnesiation (S 1 ), it can be noticed that the Mg 2+ ions interact with S atoms and show stick sediment distribution. The stick sediments disappeared in de-magnesiation state (S 2 ), and the surface showed smooth morphology. Fig. 6d and insets of Fig. 6e-g show EDS and elemental mapping results for S 0 , S 1 , and S 2 , respectively. By following the evolution of Mg 2+ ratio that increased from zero in the pristine state to 16.32 wt% upon discharge and decreased to 11.48 wt% upon recharge state, the results support the partially reversible conversion reaction of the S 8 to polysulde molecules and vice versa. Furthermore, the evolution of the sulfur ratio from 18 wt% in S 2 aer discharging to 6 wt% upon subsequent charge process conrms the shuttling of polysulde, which is presumably responsible for the capacity fading.

Conclusion
In this work, we investigated the efficiency and electrochemical compatibility of EC additive to halogen-free electrolyte based Mg(NO 3 ) 2 in acetonitrile (ACN) and tetraethylene glycol dimethyl ether (G4). The FT-IR spectra conrm that EC additive changed the Mg 2+ solvation structure via supporting weak solvent-cation interaction, promoting the rapid desolvation of Mg 2+ ions. Furthermore, the addition of EC enhanced the ionic conductivity and Mg 2+ ion transference number of the electrolyte, which can be understood from the role of EC in disassociating the magnesium salt and forming SEI. However, the introduction of EC results in the formation of a resistive SEI, causing high Mg stripping/plating overpotential. The Mg-S full cell using modied HFE (HFE_EC) delivers an initial specic capacity of 900 mA g −1 with a cycle life of up to 10 cycles in case of activating the cell with electrochemical conditioning. Much work must be dedicated to optimizing the EC-based halogenfree electrolytes to realize practical Mg-S batteries.

Conflicts of interest
There are no conicts to declare.